Heterostructured thin-film catalysts comprising nanocavities

ABSTRACT

A heterostructured catalyst includes a 2-dimensional (2D) array of titanium including nanocavities that are all directly attached to a substrate. Each of the titanium including nanocavities have a pore with a nanopore size and a wall with a nanowall thickness. The titanium including nanocavities can be titania nanocavities with a metal layer or a metal compound layer on the titania nanocavities including inside the pores, or the titanium including nanocavities can include SrTiO 3  or consist of SrTiO 3 , each with a surface layer of reduced SrTiO 3 .

FIELD

Disclosed embodiments relate to heterostructured thin film catalyststhat include cavities.

BACKGROUND

A significant amount of research recently has focused on solar energyconversion to provide clean combustible chemical fuels (H₂, CH₄, etc.)to address a future impending global energy crisis as well as hazardousenvironmental pollution conditions largely induced by fossil fuelconsumption. Solar-to-fuel conversion significantly depends onsemiconductor materials that can harvest photon energy across the widesolar spectrum (from ultraviolet (UV) to near-infrared (NIR) region) andsimultaneously generate charge-carriers on the suitable energy levelsfor H⁺ reduction. Among various semiconductor materials andtechnologies, low-cost, nontoxic, and chemical stable TiO₂ (titania) haslong been studied as a promising photocatalyst for solar-driven watersplitting. However, the main disadvantages of a wide band gap andsluggish charge transfer kinetics of TiO₂ has limited its feasibilitywithin the visible light region, which accounts for approximately 43% ofthe power output of the solar spectrum.

It is possible to generally overcome these issues with using TiO₂ by avariety of techniques, including heavily doping (e.g., doping withnitrogen), integrating narrow band gap semiconductors, and decoratingwith noble metals and co-catalysts (CdS, CdSe, Pt, Au, Ag), to broadenthe passive oxide semiconductor absorption range to include visible ornear infrared (NIR), creating a more efficient photocatalysts. However,the widespread use of scarce noble metals is not considered a reasonableoption due to their high cost and potential to be environmentally toxic.

SUMMARY

This Summary is provided to introduce a brief selection of disclosedconcepts in a simplified form that are further described below in theDetailed Description including the drawings provided. This Summary isnot intended to limit the claimed subject matter's scope.

Disclosed aspects include a heterostructured catalyst comprising a2-dimensional (2D) array of titanium comprising nanocavities that areall directly attached to a substrate, wherein each of the nanocavitieshave a pore with a nanopore size and a wall with a nanowall thickness.The nanocavities can comprise TiO₂ nanocavities and there is a metallayer or a metal compound layer on the titania nanocavities, or thenanocavities can consist of SrTiO₃. The metal layer or a metal compoundlayer on the titania nanocavities can comprise MoS₂, SrTiO₃ or a varietyof different metals or metal alloys. The 2D array of titanium comprisingnanocavities is generally a periodically ordered array having aninterval of the periodicity (center-to-center) of 10 nm to 100 nm with adistribution all within ±1 nm of a mean interval value, the nanoporesize is 10 to 200 nm with a distribution all with ±1 nm of a mean poresize value, and the nanowall thickness is 5 nm to 20 nm with adistribution all within ±0.5 nm of a mean wall thickness value.

Disclosed a heterostructured catalysts have applications including forsolar cells and for sunlight-driven hydrogen (H₂) production from water(e.g., seawater) which is considered an ultimate solution to the energyand environmental crisis. Solar fuel cells including disclosedheterostructured catalysts are also disclosed.

Disclosed aspects also include a method for forming a thin filmheterostructure catalyst including anodizing to oxidize a titaniumsurface on a substrate to form a 2D array of titania nanocavities, andthen forming a metal layer or a metal compound layer on the titaniananocavities, or converting the titania nanocavities into SrTiO₃nanocavities.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a perspective depiction of an example heterostructuredcatalyst comprising a 2D array of titania nanocavities that are alldirectly attached to a substrate, wherein each of the nanocavities havea pore with a nanopore size and a wall with a nanowall thickness, andwhere there is a metal layer or a metal compound layer on the titaniananocavities. A blow-up depiction of one the nanocavities is alsoprovided that shows the nanopore size, the nanowall thickness, and themetal layer thereon.

FIGS. 2A-D depicts the in-process structures corresponding to steps forforming an example heterostructured catalyst where the titaniananocavities are converted to SrTiO₃ nanocavities, where the surface ofthe SrTiO₃ is then reduced.

FIGS. 3A-D depicts the in-process structures corresponding to steps forforming an example heterostructured catalyst comprising forming a 2Darray of titania nanocavities that are all directly attached to asubstrate, where the metal layer or a metal compound layer on thetitania nanocavities is described as comprising aluminum as an examplemetal.

FIG. 4 depicts a solar fuel cell that includes a solar H₂ and O₂generation chamber in the middle coupled with a H₂ fuel cell on the leftand an O₂ battery cell on the right.

FIG. 5 depicts a solar cell that includes a disclosed heterostructuredcatalyst.

FIGS. 6A-G include FIG. 6A and FIG. 6B which are scanned SEM images ofTiO₂ and r-SrTiO₃, respectively. FIGS. 6C and 6D are cross-sectionalscanned SEM images of TiO₂ and r-SrTiO₃, respectively. FIG. 6E is ascanned TEM image of r-SrTiO₃. FIG. 6F is a scanned STEM image andcorresponding elemental mapping of r-SrTiO₃ (left to right: Ti, Sr, O).FIG. 6G is a scanned HRTEM image of r-SrTiO₃. The scale bars in FIGS. 6Aand 6B are 200 nm, and the scale bars in FIGS. 6C, 6D, and 6F are 100nm, in FIG. 6E the scale bar is 50 nm, and in FIG. 6G the scale bar is 5nm.

FIGS. 7A-D include FIG. 7A comprising a visible light absorption spectraof r-SrTiO₃ after exposure to oxidation in air at 400° C. over time inambient conditions. FIG. 7B is a visible light absorption spectra of theoriginal r-SrTiO₃ and recovered r-SrTiO₃. The recovered r-SrTiO₃ wasobtained by NaBH₄ reduction of the oxidized r-SrTiO₃. FIG. 7C showsincident angle-dependent light absorption spectra of r-SrTiO₃. FIG. 7Dis a schematic illustration of the incident angle-dependent lightabsorption measurement.

FIGS. 8A-E include FIG. 8A being a linear sweeps voltammogram (LSV) ofSrTiO₃ and r-SrTiO₃ photoelectrodes. FIG. 8B, and FIG. 8C are transientphotocurrent responses of SrTiO₃ and r-SrTiO₃ photoelectrodes at 1.23 Vvs. RHE under AM 1.5G illumination without and with a L-42 cutoff filter(λ≥420 nm), respectively. FIG. 8D shows IPCE spectra in the region of430-730 nm at 1.23 V vs. RHE. FIG. 8E is a schematic illustration of adisclosed mechanism of r-SrTiO₃ for photoelectrochemical (PEC) watersplitting under visible light.

FIGS. 9A-E include FIG. 9A and FIG. 9B showing scanned SEM and TEMimages of 20 nm thick aluminum on TiO₂ nanocavities represented hereinas Al20@TiO₂, respectively. FIG. 9C is a scanned HRTEM image ofAl20@TiO₂. FIG. 9D is a cross sectional STEM image of Al20@TiO₂ (wherethe inset shows the linear energy dispersive spectroscopy (EDS) profilecaptured from the dashed line, and the Al NPs are highlighted by dashedcircles). FIG. 9E is a scanned STEM image and the corresponding EDSmapping of Al20@TiO₂.

FIGS. 10A-F include FIG. 10A, and FIG. 10B comprising an XRD and XPS ofAl10@TiO₂, Al20@TiO₂, and 30 nm thick aluminum on TiO₂ nanocavitiesrepresented as Al30@TiO₂. FIG. 10C is a high-resolution XPS of Al in theAl20@TiO₂. FIG. 10D is a UV-vis absorption spectra of TiO₂, Al10@TiO₂,Al20@TiO₂, and Al30@TiO₂. FIG. 10E is an incident angle-dependent lightabsorption spectra of Al20@TiO₂. FIG. 10F is a finite difference timedomain FDTD simulation of Al20@TiO₂ as a function of incident angle.

FIGS. 11A-F include FIG. 11A comprising a linear sweep voltammograms ofTiO₂, Al10@TiO₂, Al20@TiO₂, and Al30@TiO₂. Transient photocurrentresponses of TiO₂, Al10@TiO₂, Al20@TiO₂, and Al30@TiO₂ NCAs at 1.23 Vvs. RHE under UV-visible in FIG. 11B and visible light illumination inFIG. 11C, respectively. FIG. 11D shows the photoconversion efficiency asa function of external potential. FIGS. 11E, 11F show IPCE andelectrochemical impedance spectroscopy of TiO₂, Al10 @ TiO₂, Al20 @TiO₂, and Al30 @ TiO₂.

FIGS. 12A-G include FIG. 12A comprising a SEM image of as-anodized TiO₂nanocavity arrays and FIG. 12B after e-beam deposition of 30 nm Mo; FIG.12C is a cross-sectional view of 30 nm thick MoS₂ on TiO₂ nanocavitiesrepresented as MoS₂(30)@TiO₂ (the number after MoS₂ used herein is itsthickness in nm); FIG. 12D shows 10 nm Mo sulfurized for 360 min on TiO₂nanocavity; FIG. 12E shows 10 nm Mo sulfurized for 10 min on a Ti foil;FIG. 12F shows a Mo foil sulfurized for 10 min (Scale bars: 200 nm); andFIG. 12G shows a scanned photograph and morphologies of MoS₂(10)@TiO₂ atdifferent sample locations.

FIGS. 13A-C include FIG. 13A comprising an excitation angle-resolvedplasmonic properties of MoS₂@TiO₂ (Schematic design of the experimentalsetup is inserted. The substrate can be rotated clockwise around thez-axis by an angle Φ). FIG. 13B is a UV-vis absorption spectra of 30 nmthick MoS₂, referred to as MoS₂(30). TiO₂ sulfurized at 400° C. for 10,30, 50 min which shows that the photocatalyst displays weakeningabsorption in the visible-light region with increasing sulfurizationtime, which is caused by the complementary of S in MoS₂ crystal andreducing S defects. FIG. 12C is a SERS spectra of 4-MBA recorded fromdifferent MoS₂@TiO₂ systems.

FIGS. 14A-E include FIG. 14A comprising a H₂ evolution rate overMoS₂@TiO₂ films to MoS₂ mass loading (Photocatalytic activity toseawater is inserted following. All testing carried out under simulatedsolar light). FIGS. 14B, and 14C show recycling photocatalytic H₂evolution test of MoS₂(30)@TiO₂(compact) and MoS₂@Mo(compact) over 15 h.FIG. 14D and FIG. 14E show a photoluminescence (PL) spectra of MoS₂@TiO₂cocatalysts at steady state and time-resolved transient PL decay, withcharge-carriers lifetime, inserted accordingly.

DETAILED DESCRIPTION

Disclosed embodiments are described with reference to the attachedfigures, wherein like reference numerals, are used throughout thefigures to designate similar or equivalent elements. The figures are notdrawn to scale and they are provided merely to illustrate aspectsdisclosed herein. Several disclosed aspects are described below withreference to example applications for illustration. It should beunderstood that numerous specific details, relationships, and methodsare set forth to provide a full understanding of the embodimentsdisclosed herein.

One having ordinary skill in the relevant art, however, will readilyrecognize that the disclosed embodiments can be practiced without one ormore of the specific details or with other methods. In other instances,well-known structures or operations are not shown in detail to avoidobscuring aspects disclosed herein. Disclosed embodiments are notlimited by the illustrated ordering of acts or events, as some acts mayoccur in different orders and/or concurrently with other acts or events.Furthermore, not all illustrated acts or events are required toimplement a methodology in accordance with this Disclosure.

Notwithstanding that the numerical ranges and parameters setting forththe broad scope of this Disclosure are approximations, the numericalvalues set forth in the specific examples are reported as precisely aspossible. Any numerical value, however, inherently contains certainerrors necessarily resulting from the standard deviation found in theirrespective testing measurements. Moreover, all ranges disclosed hereinare to be understood to encompass any and all sub-ranges subsumedtherein. For example, a range of “less than 10” can include any and allsub-ranges between (and including) the minimum value of zero and themaximum value of 10, that is, any and all sub-ranges having a minimumvalue of equal to or greater than zero and a maximum value of equal toor less than 10, e.g., 1 to 5.

FIG. 1 is a perspective depiction of an example heterostructuredcatalyst 100 comprising a 2D array of titania nanocavities 110 that areall directly attached to a substrate 105, wherein each of thenanocavities 110 have a pore with a nanopore size and a wall with ananowall thickness. The substrate 105 can comprise a silicon wafer, aglass wafer, or a conducting polymer.

There is a metal layer or a metal compound layer 110 c on the titaniananocavities 110. A blow-up depiction of one the nanocavities 110 isalso provided that shows dimensions representing the nanopore size, thenanowall thickness, and the metal layer 110 c on the titaniananocavities 110, which shows the metal layer 110 c being on the outsidesurface and the inside surface of titania nanocavity 110.

The metal layer or a metal compound layer 110 c on the titaniananocavities 110 can comprise metal compounds such as MoS₂, or SrTiO₃,or a variety of metals or metal alloys such as Al, Cu, Au, Ag, Pt, ortheir alloys. The 2D array of titania nanocavities 110 are shown asbeing hollow cylindrical shaped cells that are generally positioned likea crystal lattice including a highly periodically ordered array. Theinterval of the periodicity is generally 10 nm to 100 nm with adistribution all within ±1 nm of a mean interval value, the nanoporesize is generally 10 nm to 200 nm with a distribution all with ±1 nm ofa mean pore size value, and the nanowall thickness is generally 5 nm to20 nm such as 10 nm to 15 nm with a distribution all within ±0.5 nm of amean wall thickness value.

It has been found that disclosed heterostructured catalysts exhibitenhanced light absorption in a broad range of the solar spectrum, suchas from the ultraviolet (UV) light region to the visible and near-IRlight regions (generally in a range of 300 nm to 700 nm) which may beinduced by a mechanism comprising the plasmonic effect and intrinsicoptical resonance of the respective materials. The mechanism disclosedherein believed to be accurate is however not needed to practicedisclosed aspects is that disclosed metal compound nanocavities, such ascomprising TiO₂ or SrTiO₃, can absorb UV-light, the heterostructuremetal or metal compound (e.g., Al or reduced SrTiO₃) can absorb visiblelight, and metal compounds such as MoS₂ or Ni(OH)₂ can absorb visibleand near-IR light.

FIGS. 2A-D depict the in-process structures corresponding to steps foran example heterostructured catalyst where the titania nanocavities areconverted to SrTiO₃ nanocavities, which are then reduced at the surfaceto form r-SrTiO₃ surface. This embodiment recognizes plasmon resonance(LSPR) have a high potential for solar energy harvesting applications.Exploring nonmetallic plasmonic materials is desirable yet challenging.Herein, an efficient nonmetallic plasmonic perovskite photoelectrode,namely, SrTiO₃, with a periodically ordered 2D nanoporous structureshowing an intense LSPR in the visible light region is disclosed. Thecrystalline-core@amorphous-shell structure of the SrTiO₃ photoelectrodeenables a strong LSPR due to the high charge carrier density induced byoxygen vacancies in the amorphous shell. The reversible tunability inLSPR of the SrTiO₃ photoelectrode was observed by oxidation/reductiontreatment and incident angle adjusting. Such a nonmetallic plasmonicSrTiO₃ photoelectrode displays a large plasmon-enhancedphotoelectrochemical water splitting performance with a photocurrentdensity of 170.0 μA cm⁻² under visible light illumination and a maximumincident photon-to-current-conversion efficiency of 4.0% in the visiblelight region, which are comparable to the state-of-the-art plasmonicnoble metal sensitized photoelectrodes. FIGS. 2A-2D are described infurther detail below.

FIGS. 3A-D depict the in-process structures corresponding to steps informing a disclosed heterostructure catalyst, where the metal layer or ametal compound layer on the titania nanocavities comprises aluminumwhich is used only as an example of a metal or metal alloy coating. Thisembodiment comprises Al on titiania naoncavities (abbreviated asAl@TiO₂) heterostructures for efficient PEC water splitting bycontrollably isolating Aluminum (Al) nanoparticles (NPs) individuallyinto disclosed TiO₂ nanocavity arrays (NCAs). Compared with bare TiO₂,Al@TiO₂ shows enhanced PEC performance under solar light illumination.The significantly enhanced PEC activity of an Al@TiO₂ photoanode isattributed to the LSPR induced electromagnetic field enhancement inUV-visible region and photoinduced hot carrier of Al NPs. Moreover, thenaturally formed Al₂O₃ thin surface layer formed in air at roomtemperature sometimes termed a native oxide can act as a protectivelayer by restraining Al NPs corrosion and reducing the surface chargerecombination. This aspect evidences that a novel plasmonic system canbe rationalized by forming isolated earth-abundant metal into thenanostructured semiconductor film with a highly ordered morphology,which is believed to open a new paradigm for an advanced catalyst designfor PEC water splitting. FIGS. 3A-3D are described in further detailbelow.

FIG. 4 depicts a solar fuel cell 400 that includes a disclosedheterostructured catalyst shown as 100 on an anode of a PV assembly 410comprising solar cell(s) including an active region comprising a p-typematerial and an n-type material positioned between the anode and acathode, where the heterostructured catalyst 100 in operation is exposedto sunlight 450. The substrate of the heterostructured catalyst shown as100 is electrically conductive and is used to make a connection with theexternal circuit, acting like a current collector. The solar fuel cell400 converts light from the sun shown as 450 to run the electron andproton flow in reverse as compared to a conventional solar cell or solarmodule inclising a plurality of solar cells. Coupling the photogeneratedelectrons and protons from the PV cell(s) to the heterostructuredcatalyst 100 breaks the OH bonds of water (H₂O) supplied which generatesH₂ and O₂. The H₂ is shown used as a chemical fuel in the H₂ fuel cell420. O₂ is used as a chemical fuel shown used by an O₂ battery cell 430(e.g., a lithium-O₂ battery). The H₂ fuel cell 420 and O₂ battery cell430 each are shown generating electricity both shown providingelectricity out.

FIG. 5 depicts a solar cell shown by example as a solar dye cell 500that includes a disclosed heterostructured catalyst 100. The electricalpower generated by the solar dye cell 500 is shown as a voltagerepresented as V″. The heterostructured catalyst 100 is positionedbetween the dye with electrolyte 510 and the bottom glass/Indium tinoxide (ITO) layer 520 that generally has a Pt or a graphite coating.Above the dye with electrolyte 510 is a top glass/ITO layer 540 thatgenerally also has a Pt or a graphite coating. The glass/ITO 540 thatgenerally has a Pt or a graphite coating functions as a counterelectrode and the heterostructured catalyst 100 functions as a workingelectrode, which are separated from one another by the dye withelectrolyte 510. No membrane is needed by the solar dye cell 500.

A disclosed aspect with a metal compound on the nanocavities or beingthe entire nanocavities is now described as being SrTiO₃. H₂ productionvia PEC splitting of water has been developed as a technology for solarenergy harvesting. The development of efficient photoanodes is neededfor practical PEC applications. However, most recently developedphotoanodes still suffer from unsatisfactory solar energy conversionefficiencies, largely due to their insufficient light absorption, lowcharge separation, and low transfer efficiencies. One effective strategyfor expanding the light absorption and improving the efficiencies is tocombine the photoanodes with plasmonic noble metals due to the LSPR. Forexample, dramatically plasmon-enhanced PEC water splitting performanceshave been achieved over some photoanodes, such as TiO₂, BiVO₄, and ZnO,via decorating them with plasmonic Au nanostructures. However, theplasmonic metals, especially noble metals, are rare and are expensive,which limit their wide applications.

A disclosed aspect is a plasmonic SrTiO₃ photoelectrode with aperiodically ordered nanoporous structure formed by hydrothermallyconverting a TiO₂ nanoporous film comprising a 2D array of nanocavities,that is followed by a reduction process. The as-prepared surface reducedSrTiO₃ (r-SrTiO₃) that has a crystalline-core@amorphous-shell structurewith abundant oxygen vacancies in the amorphous shell can absorb visiblelight due to the LSPR property. The reversible tunability in theplasmonic resonance of r-SrTiO₃ can be observed through a simpleoxidation/reduction process and adjusting the incident angles. Such anr-SrTiO₃ photoelectrode exhibits a dramatically enhanced PEC watersplitting performance compared with a pristine SrTiO₃ photoelectrodeunder visible light irradiation.

An example synthetic procedure for a plasmonic r-SrTiO₃ photoelectrodebegins in FIG. 2A that shows a Ti foil. A highly ordered TiO₂ nanoporousfilm is first obtained by Ti anodization to form a TiO₂ nanoporous filmwith the results shown in FIG. 2B. Example anodization processparameters are described in the Examples section below. A SEM imageobtained revealed that the as-obtained TiO₂ nanoporous film shows ahoneycomb-shaped morphology with a nanopore size of about 60 nm and filmthickness of 120 nm. Then the as-obtained TiO₂ nanoporous film washydrothermally converted into a SrTiO₃ nanoporous film followed by anannealing treatment with the result shown in FIG. 2C. This 120 nm thickTiO₂ nanoporous film was chosen because it can be completely convertedto SrTiO₃ without destroying the porous structure. The obtained SrTiO₃film shows a highly porous structure, which is very similar to theannealed TiO₂ nanoporous film.

However, the nanocavity wall of SrTiO₃ (˜20 nm) generally becomesthicker as compared with the TiO₂ nanoporous film (˜10 nm) due to thephase transition and lattice expansion. Finally, the plasmonic r-SrTiO₃nanoporous film was obtained by reducing the obtained SrTiO₃ nanoporousfilm with NaBH₄ in a N₂ atmosphere, with the result shown in FIG. 2D.r-SrTiO₃ on the surface of SrTiO₃ provides a heterostructure that isneeded to generate the plasmonic effect. The r-SrTiO₃ nanoporous filmshows almost the same morphology as the SrTiO₃ nanoporous film. an EDSanalysis indicated that only Ti, O, and Sr elements are detected fromthe r-SrTiO₃ nanoporous film. TEM of the r-SrTiO₃ nanoporous film showeda honeycomb-shaped morphology, which is consistent with the SEM. Thehigh-angle annular-dark-field (HAADF) scanning transmission electronmicroscopy (STEM) image and the corresponding elemental mapping ofr-SrTiO₃ show a uniform distribution of Ti, Sr, and O, furtherconfirming the formation of SrTiO₃.

FIGS. 6A-G include 6A and FIG. 6B are scanned SEM images of TiO2 andr-SrTiO3, respectively. FIGS. 6C and 6D are cross-sectional scanned SEMimages of TiO2 and r-SrTiO3, respectively. FIG. 6E is a scanned TEMimage of r-SrTiO3. FIG. 6F is a scanned STEM image and correspondingelemental mapping of r-SrTiO3 (left to right: Ti, Sr, O). The dashedparts in 6G) indicate nanocavities. FIG. 6G is a scanned HRTEM image ofr-SrTiO3. The Scale bars in 6 a and 6B are 200 nm, and in 6C, D, and F100 nm, FIG. 6E 50 nm, and in FIG. 6G the scale bars are 5 nm.

A high-resolution TEM (HRTEM) of r-SrTiO₃ shown in FIG. 6G evidencesthat an amorphous layer that is 4 nm thick which is apparently formed onthe surface. Moreover, the core of r-SrTiO₃ is still highly crystallinewith a measured lattice spacing of 0.27 nm, which is consistent with thecubic SrTiO₃ (110) plane. The HRTEM obtained confirmed that the r-SrTiO₃nanoporous film has a crystalline-core@amorphous-shell structure. Duringthe preparation of r-SrTiO₃, NaBH₄ was used as an efficient oxygenscavenger by thermally decomposing and generating active hydrogen. Theactive hydrogen has a strong reductive capability, which can react withthe oxygen atoms on the surface of SrTiO₃, thus producing a disordered(oxygen vacancies) surface layer. As a comparison, the pristine SrTiO₃without reducing treatment is highly crystallized, as observed from theclear lattice feature displayed in a HRTEM image obtained.

The crystalline phases of the materials were recorded by X-raydiffraction (XRD). In contrast to the previous reports, which usuallycontained incompletely converted TiO₂ via the similar hydrothermalmethod by using micrometer-sized TiO₂ nanotubes/nanorods as precursors,the obtained SrTiO₃ nanoporous film showed excellent phase purity,indicating a complete phase conversion. The properties of the TiO₂ film,such as porous structure, ultrathin cavity wall (˜10 nm), and thin filmthickness (˜120 nm), are beneficial for the complete conversion toSrTiO₃. The diffraction peaks are ascribed to metallic Ti and cubicSrTiO₃ phase and no XRD peaks of TiO₂ or other composition were found.

The purity of the SrTiO₃ nanoporous film was further confirmed by Ramanspectra. The SrTiO₃ demonstrated two main regions of Raman peaks at250-400 and 600-750 cm, which are in line with the standard Ramanspectrum of SrTiO₃ As for the r-SrTiO₃, the XRD peaks and Raman peaksexhibit slight broadening and intensity decreasing compared with thepristine SrTiO3, which are due to the existence of oxygen vacancies. Incontrast, direct annealing of the TiO₂ nanoporous film at 450° C. onlyleads to the production of crystalline TiO₂, which is revealed by theXRD and Raman spectra.

FIGS. 7A-D include FIG. 7A comprising a visible light absorption spectraof r-SrTiO3 after exposure to oxidation in air at 400° C. over time inambient conditions. FIG. 7B is a visible light absorption spectra of theoriginal r-SrTiO3 and recovered r-SrTiO3. The recovered r-SrTiO3 wasobtained by NaBH4 reduction of the oxidized r-SrTiO3. 7C shows incidentangle-dependent light absorption spectra of r-SrTiO3. FIG. 7D is aschematic illustration of the incident angle-dependent light absorptionmeasurement.

FIGS. 8A-E include FIG. 8A being a linear sweeps voltammogram (LSV) ofSrTiO₃ and r-SrTiO₃ photoelectrodes. FIGS. 8B, and C are transientphotocurrent responses of SrTiO₃ and r-SrTiO₃ photoelectrodes at 1.23 Vvs. RHE under AM 1.5G illumination without and with a L-42 cutoff filter(2>420 nm), re-spectively. FIG. 8D shows IPCE spectra in the region of430-730 nm at 1.23 V vs. RHE. FIG. 8E is a schematic illustration of adisclosed mechanism of r-SrTiO₃ for PEC water splitting under visiblelight.

A disclosed aspect with a metal or metal alloy on the nanocavities isnow described as being aluminium. Among the earth abundant non-noblemetals, Al has been reported as an economic plasmonic alternative tonoble ones. Aluminum allows tuning the LSPR energy from the deepultraviolet to the infrared region by changing the particle size. As aconsequence, compared with noble metal-based plasmonic materials,Al-based plasmonic nanostructure should be considered for a plasmonicphotoelectrode in the scalable application of PEC water splitting.However, due to the poor dewetting property of the aluminum and thenaturally formed alumina (Al₂O₃) layer, the traditional dielectricsubstrates with a flat or disordered surface such as Si wafer andcompact dielectric films, greatly hinder the formation of periodicallyordered Al NPs for PEC applications. In addition, the formation andutilization of plasmonic Al@TiO₂ nanocavity arrays (NCAs) with a highlyordered morphology is not yet well examined.

Al@TiO₂ plasmonic heterostructure can be realized depositing theisolated Al NPs into the TiO₂ NCAs as a photoanode for PEC watersplitting. In order to examine the formation of Al/TiO₂ NCAs and theircorresponding solar absorption and charge transfer kinetics, differentthickness (10 nm, 20 nm, and 30 nm) of Al film were deposited onto theTiO₂ NCAs followed by thermal dewetting treatment. The achievedphotocurrent of Al@TiO₂ is much higher than that of the bare TiO₂ NCAs.The significantly enhanced PEC performance of Al@TiO₂ was attributed tothe preferable plasmonic heterostructure with predominant hot carriertransfer and efficient utilization of solar light.

The synthesis procedure of Al@TiO₂ is demonstrated beginning on FIG. 3Awith a Ti foil. Firstly, anodization treatment is used to form highlyordered TiO₂ NCAs with porous nanostructure as shown in FIG. 3B with ananopore size of 50 nm and a nanowall thickness of 8 nm. Then, TiO₂ NCAswere deposited with Al thin layers with different thickness (10 nm, 20nm, and 30 nm) by e-beam evaporation with the result shown in FIG. 3C.Eventually, the isolated Al NPs were deposited into the TiO₂ NCAs by athermal dewetting treatment in a nitrogen atmosphere with a schematicimage of the Al@TiO₂ nancavities shown in FIG. 3D, which was thendirectly used as photoanodes without further treatment for PEC watersplitting.

FIGS. 9A-F include FIG. 9A and FIG. 9B showing SEM and TEM images ofAl20@TiO₂. FIG. 9C is a HRTEM image of Al20@TiO₂. FIG. 9D is a crosssectional STEM image of Al20@TiO₂ (inset shows the linear EDS profilecaptured from the dashed line. The Al NPs are highlighted by whitedashed circles). FIG. 9E is a STEM image and the corresponding EDSmapping of Al20@TiO₂.

FIGS. 10A-F include FIG. 10A, and FIG. 10B comprising XRD and XPS ofAl10@TiO₂, Al20@TiO₂, and Al30@TiO₂, which as described above mean 10,20, and 30 nm thick aluminum on TiO₂ nanocavities. FIG. 10C is ahigh-resolution XPS of Al in the Al20@TiO₂. 10D is a UV-vis absorptionspectra of TiO₂, Al10@TiO₂, Al20@TiO₂, and Al30@TiO₂. 10E is an incidentangle-dependent light absorption spectra of Al20@TiO₂. FIG. 10F is aFDTD simulation of Al20@TiO₂ as a function of incident angle.

FIGS. 11A-F include FIG. 11A comprising a linear sweep voltammograms ofTiO₂, Al10@TiO₂, Al20@TiO₂, and Al30@TiO₂. Transient photocurrentresponses of TiO₂, Al10@TiO₂, Al20@TiO₂, and Al30@TiO₂ NCAs at 1.23 Vvs. RHE under UV-visible in FIG. 11B and visible light illumination inFIG. 11C, respectively. FIG. 11D shows the photoconversion efficiency asa function of external potential. FIGS. 11E, and 11F show IPCE andelectrochemical impedance spectroscopy of TiO₂, Al10@TiO₂, Al20@TiO₂,and Al30@TiO₂. SEM and transmission electron microscopy (TEM) wereemployed to investigate the influence of the deposited Al film thicknesson the thermal dewetting kinetics of Al onto the nanostructured TiO₂NCAs. The Al₁₀@TiO₂ NCAs (10 nm Al layer) showed that the averagediameter of Al NPs inside the pores of TiO₂ NCAs was 25 nm and meanwhilea “crown” nanostructure was formed at the edges of TiO₂ NCAs wall as aresult of the fewer nucleation centers created by 10 nm Al layer. Withthe increase of the Al layer thickness up to 20 nm (Al₂₀@TiO₂), theperiodically ordered Al NPs with the diameter of 34 nm were uniformlyisolated inside each pore of TiO₂ NCAs.

However, further increasing the Al layer thickness to 30 nm (Al₃₀@TiO₂)will generate the over-grown Al NPs with the size of 42 nm inside of theTiO₂ NCAs. At the same time, larger Al NPs with the size of 70 nm wereaggregated and nonuniformly coated on the surface of TiO₂ NCAsindicating the saturation of Al NPs loading. The over-grown Al NPs onthe TiO₂ NCAs would reduce the light-receiving area of TiO₂ NCAs andblock the open pores of TiO₂ NCAs, leading to a negative adverse effecton the PEC performance. High-resolution TEM (HRTEM) and scanningtransmission electron microscopy (STEM) images further confirmed theformation of the heterostructured Al₂₀@TiO₂. The lattice fringe spacingsof 0.23 nm and 0.35 nm correspond to the (111) plane of Al and (101)plane of anatase TiO₂, respectively.

Moreover, an amorphous Al₂O₃ ultra-thin layer (1.5 nm thick) wasconformally/naturally formed on the surface of Al NPs, which played apositive role in preventing Al NPs from chemical corrosion, contributingto a reduced recombination rate and eventually facilitating performanceof PEC water splitting. The cross-sectional TEM view and linear energydispersive spectroscopy (EDS) profile obtained showed that the highlydiscrete Al NPs were formed inside the TiO₂ NCAs. TEM-EDS mapping showedthe distribution of Ti, O, and Al, indicating that each Al NP weresolely isolated in the individual pore of TiO₂ NCAs.

The crystal structures of Al@TiO₂ NCAs were characterized by X-raydiffraction measurements. The main diffraction peaks located at 35.06°,38.43°, 40.23°, and 53.04° correspond to Ti substrate. Thecharacteristic peak at 25.3° is ascribed to the (101) crystal plane ofanatase TiO₂. However, the characteristic peaks ascribed to Al isinvisible in XRD patterns due to the ultra-low content of Al, which isoverlapped the by Ti peak at 38°. X-ray photoelectron spectroscopy wasemployed to investigate the chemical composition of Al@TiO₂ NCAs,revealing the existence of Ti, O, and Al. High-resolution XPS spectra ofTi 2p revealed that two peaks at 459.20 eV and 464.97 eV were ascribedto Ti 2p_(3/2) and Ti 2p_(1/2) in TiO₂. Basically, an ultra-thin Al₂O₃layer would be naturally formed on the surface of Al NPs right afterexposure to air. Therefore, Al 2p peaks were resolved into two peaks at74.40 eV and 72.84 eV, which were attributed to Al 2p_(3/2) in Al₂O₃ andmetallic Al, respectively.

The Al/Al₂O₃ ratio was estimated from the high-resolution XPS spectra ofAl 2p, revealing that Al₃₀@TiO₂ was of the lowest Al₂O₃ content becauseof the less oxidation activity of Al NPs with larger diameter comparedwith that of Al₁₀@TiO₂ and Al₂₀@TiO₂. This result is strongly supportedby the Cabrera-Mott model that proves smaller particles are much easierto be oxidized. On the other hand, as indicated by HRTEM, the thicknessof the amorphous Al₂O₃ ultra-thin layer was approximately 2.5 nm(Al₁₀@TiO₂), 1.5 nm (Al₂₀@TiO₂), and 0.8 nm (Al₃₀@TiO₂). Note thatsuitable thickness of Al₂O₃ will not only protect Al NPs from furtheroxidation but also prevent charge recombination and eliminate theForster energy transfer. However, if the Al₂O₃ layer is thicker than theelectron tunneling limit, it will generally be adverse to interfacialcharge transfer. Eventually, a thicker Al₂O₃ layer of Al₁₀@TiO₂ willshow inferior PEC performance compared with Al₂₀@TiO₂ and Al₃₀@TiO₂.

In order to investigate the optical properties of Al@TiO₂ NCAs, theUV-visible absorption spectra were examined. The optical absorption ofTiO₂ NCAs decreased significantly at a wavelength of 380 nm, owing tothe intrinsic light absorption of anatase TiO₂. In addition, plasmonicabsorption peak centered at 462 nm, 560 nm, and 579 nm were observed inAl₁₀@TiO₂, Al₂₀@TiO₂, and Al₃₀@TiO₂, respectively, due to the LSPReffect and electromagnetic field polarization in Al@TiO₂ NCAs. On theother hand, it is noteworthy that Al₂₀@TiO₂ possesses the strongestintensity of the plasmonic absorption peak which is attributed to theoptimal plasmonic heterostructure with a narrow size distribution of AlNPs inside TiO₂ NCAs without any interference from smaller Al NPs at thetubular edge of TiO₂ NCAs. Consequently, the unique optical property ofAl₂₀@TiO₂ plasmonic heterostructure may contribute to preeminentcatalytic performance for PEC water splitting. Furthermore, the incidentangle-resolved plasmonic properties of Al₂₀@TiO₂ was also investigated.The Al₂₀@TiO₂ sample was rotated along the out-of-plane direction tochange the incident angle from 10° to 70° at an interval of 10°.

As the incident angle increased, the plasmonic absorption peaks showed ablue shift (to lower wavelength) which was able to return backward withthe decreases of incident angle. The similar phenomena were observed inthe Al₁₀@TiO₂ and Al₃₀@TiO₂ NCAs as well. This dynamic shift resultingfrom the incident angle dependence was further confirmed by the FDTDsimulation, showing that an approximately 60 nm blue shift occurred withincident angle increasing from 10° to 70°. It revealed that theorientation of Al NPs inside TiO₂ NCAs was very sensitive to theplasmonic resonance frequency mode induced by angle-resolved incidentlight and it had a significant impact on the distribution of carrierconcentration inside the plasmonic catalyst.

Surface-enhanced Raman scattering (SERS) and angle-resolved SERS werefurther employed to examine the localized electromagnetic fieldenhancement surrounding the Al NPs inside TiO₂ NCAs. Rhodamine B (RhB)with the concentration of 10⁻⁶ M was probed with the excitationwavelength of 785 nm for Raman measurements. All the Al@TiO₂ NCAs showedcharacteristic Raman peaks of RhB, namely, the characteristic peaks at1125 cm⁻¹ and 1194 cm⁻¹, belonging to aromatic C—H bending. The otherfour peaks located at 1364 cm⁻¹, 1507 cm⁻¹, 1534 cm⁻¹ and 1564 cm⁻¹ wereattributed to achromatic C—C stretching. According to theelectromagnetic enhancement theory, the high SERS intensity of theAl@TiO₂ NCAs can be ascribed to their strong LSPR effect. However, noobvious Raman peaks could be found in the bare TiO₂ NCAs, indicating thestrong plasmonic effect of Al NPs in the Al@TiO₂ NCAs. Furthermore,compared with Al₁₀@TiO₂ and Al₃₀@TiO₂, Al₂₀@TiO₂ showed the strongestSERS enhancement, revealing the strongest electromagnetic fieldenhancement in the rationalized Al₂₀@TiO₂ NCAs. Moreover, angle-resolvedSERS measurement on the Al₂₀@TiO₂ NCAs was performed by mapping the SERSsignal at the different incident laser angles. Because the dispersion ofplasmonic energies changes with incident angle, the SERS enhancementvaries consistently with the detection angles Specifically, the SERSperformance decreases with the increased incident angle, which agreeswell with the discussion on the incident angle-resolved plasmonicproperties.

3D FDTD simulations were carried out to reveal the distribution ofLSPR-induced electromagnetic field in the Al₂₀@TiO₂ plasmonicheterostructure and its contribution to PEC water splitting. Thesimulation model of the heterostructure was built upon the TEM, wherethe Al NPs loaded onto the tubular wall of TiO₂ NCAs. The excitationwavelengths of 380 nm and 567 nm were used in FDTD simulation. The FDTDsuggests that the localized electromagnetic field is greatly enhanced atan incident light of 567 nm due to the LSPR absorption. Meanwhile,plasmonic hot spot region is formed near the edge of Al NPs (x-z plane)and the Al/TiO₂ interface (x-y plane), which confirms an enhancement ofPEC activity under visible light. Specifically, hot electrons willtransfer to the conduction band of TiO₂, leaving energetic holes on theAl NPs for water oxidation under light illumination. By contrast, theslight increase in the electromagnetic field excited by the UV light wasobserved, which played an equally crucial role in the enhanced PECperformance due to the synergism between the increased electromagneticfield in the UV region and the TiO₂ absorption.

The PEC performance of Al@TiO₂ and bare TiO₂ NCAs were measured in astandard three-electrode system in 0.5 M Na₂SO₄ aqueous electrolyte.Linear sweep voltammograms (LSV) of Al@TiO₂ and bare TiO₂ NCAs weretested in dark and light illumination. Almost no current was detected inthe dark. When switched to light illumination, the remarkablephotocurrent was displayed on the Al@TiO₂ NCAs. Note that, Al@TiO₂ NCAsshowed substantially larger photocurrent than that of bare TiO₂ NCAs,which is almost 2 to 3 times the photocurrent of the bare TiO₂ at 1.23 Vvs. RHE. This result proved that Al NPs played a favorable role inimproving the PEC performance of TiO₂ upon illumination.

Among all of the Al@TiO₂ NCAs, Al₂₀@TiO₂ exhibited the highestphotocurrent compared with Al₁₀@TiO₂ and Al₃₀@TiO₂. It suggests thatcritical size and loading amount of Al NPs, as well as the thickness ofthe surface Al₂O₃ ultra-thin layer, play significant roles in PEC wateroxidation reaction. Confirmed by photoabsorption spectra and SERSanalysis, with the increase of Al loading amount (from 10 to 20 nm), AlNPs were uniformly distributed inside TiO₂ NCAs and thus the number ofhot electrons increased, which was favorable for PEC reactions. However,when the Al loading amount increased to 30 nm, particularly Al NPs wereaggregated on the surface of TiO₂ NCAs which reduced the light-receivingarea (shading effect) and thus affected adversely the PEC performance.On the other hand, the Al₂O₃ passivation layer is also critical to theimprovement of PEC activity. The Al₂O₃ passivation layer with a properthickness contributes to preventing Al NPs from corrosion and reducesthe surface charge recombination. However, if it exceeds the thicknesslimit for charge tunneling, for example, Al₁₀@TiO₂, charge transfer willbe blocked to across the photoelectrode/electrolyte interface. Thus, asa result of the insufficient loading amount of Al NPs and thicker Al₂O₃layer, Al₁₀@TiO₂ showed inferior PEC activity compared with Al₂₀@TiO₂and Al₃₀@TiO₂ NCAs.

The chronoamperometric I-t curves were measured to examine the instantphotoresponse of Al@TiO₂ and bare TiO₂ NCAs under the potential of 1.23V vs. RHE. The photocurrent of Al@TiO₂ sharply increases by 2-3 timesgreater than that of the bare TiO₂ NCAs under the UV-visible light whichagrees well with the LSV curves, suggesting an efficient chargeseparation and stable photocurrent in the Al@TiO₂ NCAs. The longtimephotocurrent measurement indicates that the Al₂₀@TiO₂ is stable in PECwater splitting reaction). After the longtime reaction, the Al₂₀@TiO₂was collected and investigated by SEM, which showed no obvious change ofmorphology, further confirming the good stability of Al₂₀@TiO₂ for PECwater splitting. To further confirm the contribution from the plasmoniceffect of Al NPs, I-t curves under visible-light illumination (AM 1.5 Gcombined with an L-42 cutoff filter, λ≥420 nm) were also recorded. Noobvious photocurrent can be observed from the bare TiO₂ NCAs due to thelarge bandgap and insufficient excitation under visible light.Reproducible photocurrent can be obtained over Al@TiO₂, revealing animproved catalytic performance of TiO₂ NCAs after Al implantation undervisible light, due to the LSPR effect. Consequently, the photoconversionefficiency (f) as a function of external potential was calculated byequation (1):

η=I(1.23−V _(app))/P _(light)  (1)

where V_(app) is the applied external potential vs. RHE; P_(light) isthe power density under illumination and I is the externally recordedphotocurrent density. A greatly improved photoconversion efficiency ofAl₂₀@TiO₂ was achieved compared with Al₁₀@TiO₂, Al₃₀@TiO₂, and bare TiO₂NCAs.

Furthermore, incident-photon-to-current-conversion efficiency (IPCE)tests were performed to investigate the PEC activities of Al@TiO₂ andTiO₂ NCAs. Significantly, the shape of the IPCE curve in the visibleregion is similar to that of the LSPR absorption spectra. In addition,steeply increased transient photocurrent of Al@TiO₂ under visible lightresponse is also consistent well with the IPCE of Al@TiO₂.

These results strongly validate that the LSPR of Al NPs contributes toPEC activity in visible light. On the other hand, it is worth mentioningthat the IPCE data from photoanode of Al₂₀@TiO₂ indicates a higherphotocurrent density compared with that of Al₁₀@TiO₂ and Al₃₀@TiO₂. Itis mainly due to the synergetic effect of reduced charge transferresistance and longer lifetime of hot carriers within the optimizedAl₂₀@TiO₂ photoanode. Herein, to have an insightful understanding of thecharge transfer and lifetime, electrochemical impedance spectroscopy(EIS) and hot carrier lifetime (τ_(n)) vs. open-circuit photovoltage(V_(oc)) were employed. The resistance of electrolyte-semiconductorinterface (R₂) of Al@TiO₂ photoanode is lower than that of TiO₂ NCAs,indicating that the internal electromagnetic field built up at Al/TiO₂interface leads to effective electron-hole separation and thus reducesthe chance for charge recombination. Al₂O₃ as insulating layer willmodify the distance between TiO₂ NACs and Al, and therefore eliminatingthe Forster energy transfer to reduce the recombination of chargecarriers.

Moreover, the Al₂O₃ passivation layer formed on the surface of Al NPscan effectively prevent carriers from being captured by the surfacestates of TiO₂ NACs and significantly decrease the recombination. Whencompared with Al₁₀@TiO₂ and Al₃₀@TiO₂, Al₂₀@TiO₂ with a greatly reducedresistance makes a favorable condition for hole output from photoanodeto the electrolyte, leading to much better performance for wateroxidation. Additionally, the potential decay curves obtained viainterruption of visible-light irradiation were measured to gain aninsightful understanding of kinetical charge transfer for Al@TiO₂ NCAs.The Al@TiO₂ photoanodes were firstly irradiated with visible light forseveral seconds to induce the hot electrons transfer from Al NPs toTiO₂. When the light was interrupted, potential decayed until reachingto an equilibrium state because the hot electrons on TiO₂ would transferback to the to Al NPs. The mean lifetime of hot electrons in theconduction band of TiO₂ can be calculated by equation (2):

τ_(n)=−(k _(B) T/q)(dV _(oc) /dt)⁻¹  (2)

where k_(b) is the Boltzmann constant, q is the elementary charge of theelectron, V_(oc)=E_(ph)−E_(dark). The τ_(n) of Al₂₀@TiO₂ is larger thanthat of Al₁₀@TiO₂ and Al₃₀@TiO₂ NCAs, revealing that greatly enhancedcharge separation efficiency of hot carriers by suppressing the backelectron transfer.

Based on above discussion, the electron-transfer mechanism of theplasmonic Al@TiO₂ heterostructure for PEC water oxidation under visiblelight appears likely. When the photon energy matches the LSPR region ofAl NPs inside the TiO₂ NCAs, it will excite the LSPR absorption. Thenhot electron-hole pairs generated by the surface plasmon will beseparated at Al/TiO₂ interface and transfer towards opposite direction.A large number of hot electrons will be dynamically injected into theconduction band of TiO₂. Meanwhile, the energetic holes on the Al NPswill be utilized for water oxidation. Therefore, the improved PECactivity of Al@TiO₂ NCAs in visible light is primarily attributed to thehot carriers produced by LSPR excitation of Al NPs.

Another disclosed aspect uses MoS₂ as the metal or metal compoundcoating on the nanocavities. MoS₂ belongs to the two-dimensional (2D)layered transition metal dichalcogenides (TMDs) family that has asandwich-like structure of Mo atoms between two layers of hexagonallypacked sulfur atoms. The weak Van der Waals bonding between these 2Dlayers can give rise to single- or few-layer nanosheet architectures.This aspect recognizes that MoS₂ can be a promising electrocatalyst forthe hydrogen evolution reaction (HER), owing to the nanosized MoS₂ edgedefects that are preferential to hydrogen adsorption, so thatfew-layered nanoscale MoS₂ flakes can serve to improve the efficiency ofhydrogen evolution. Chemical exfoliation and solvothermal methods areconventionally used for the fabrication of nanostructured MoS₂. However,these methods usually promote irregular particle formation orundesirable stacking of multilayer MoS₂ deposit products. Multilayerstacking exposes more catalyst basal planes than edges, rendering muchof the material catalytically inert. Additionally, without properdispersion and immobilization, powdered nanomaterials can suffer fromparticle aggregation, leading to performance degradation. Until thisDisclosure, an efficient scheme or design principle for the integrationof ordered, nanostructured MoS₂ with wide band gap semiconductingbehavior is believed to have eluded researchers.

Compared with metallic 1T-MoS₂, 2H—MoS₂ exhibits high stability andsemiconducting properties at room temperature, allowing them to be usedas a co-catalyst coupled with other wide band gap semiconductors for H₂evolution in a photoelectron reactive medium. However, MoS₂ is not knownas the main photocatalyst, especially those that attribute to NIRintrinsic absorbance. Theoretical calculations demonstrate that the bandgap of MoS₂ can be modulated through control of particle size. In thiscase, the band gap broadens from 1.2 eV to 1.9 eV when the MoS₂architecture changes from bulk to monolayer. This characteristic can beascribed to the quantum confinement of nanomaterials. The energy bandstructure can also be affected by the metal-chalcogenide stoichiometricratio. The nonstoichiometric metal-chalcogenides (e.g., WO_(3-x),Cu_(2-x)S, and MoO_(3-x)) can display evidence of an indirect plasmonicabsorption, which is distinctly different from previously reported bandgap transitions. Therefore, LSPR can be used to describe the lightharvest phenomena, which is mainly deduced from charge collectiveoscillation on the metal chalcogenide surface propagated by numerousanions (O or S) vacancies within the crystal lattice. The nonmetal MoS₂with plasmonic absorption may be a solution to fill the visible lightharvest gap vacated by wide band gap semiconductors, effectivelyimproving the solar-to-energy efficiency of HER photocatalysts. Lastly,MoS₂@TiO₂ hybrid catalysts have been recently reported that aim toenhance photocatalytic efficiency by modulating the TiO₂ energy levelwith inter-band coupling. Disclosed MoS₂@TiO₂ heterostructures haveunique feature including functioning as a thin-film catalyst, are highlyordered and well-controlled morphology, self-cleaning, and long lifespanas compared to known MoS₂@TiO₂ hybrid catalysts.

This aspect includes a combined physical vapor deposition (PVD) andchemical vapor deposition (CVD) process for coating few layered MoS₂nanoflakes conformally on the inner surface of anodized TiO₂ nanocavityarrays (referred to as MoS₂@TiO₂) with a highly-ordered 3D hierarchicalconfiguration. The stoichiometric ratio of Mo and S atoms within theMoS₂ lattice and vertically contacting facets of MoS₂ and TiO₂ can becontrolled by tuning the S source concentration in a Na₂S_(x) solutionand altering the CVD reaction rate (see the below described experimentaldetails). This highly localized growth enables conformal and uniformMoS₂ nanoflakes to be formed on the surface of TiO₂ nanocavities.

Disclosed heterostructures with MoS₂ have shown powerful photonharvesting abilities in the UV-Vis-NIR range by tethering the TiO₂substrate to a plasmonic/intrinsic MoS₂ coating. The facilitatedelectron transfer pathway and appropriately tuned the energetic positionof the conduction band results in facile charge carrier separation anddramatically enhanced H₂ evolution efficiency. UV-Vis spectroscopyanalysis, finite element method simulation (FEM), and the monochromaticlight irradiated H₂ generation rate evidence that LSPR, mainly excitedin the wavelength from 400 nm to 600 nm, substantially contributes tophotocatalytic activity. Minimal red and NIR light induced H₂ productionis observed from the lower photoelectron energetic level of the MoS₂inter-band excitation and excessive photoelectron-hole recombination.This nonmetal plasmonic heterostructure is also expected to beapplicable to other 2D material systems, which can serve as a new designprotocol for highly efficient photocatalysts.

Conventional photocatalysts can generally only absorb UV-vis light whichoccupies only less than 50% solar light spectrum. Disclosed MoS₂/TiO₂catalysts can absorb solar energy in a much broader wavelength rangecompared to conventional photocatalysts, generally resulting at leasttwo-fold higher energy conversion efficiency compared to conventionalphotocatalyst material uses including solar energy harvesting,photovoltaics, and clean fuels generation.

A broad spectral response was evidenced ranging from ultraviolet-visible(UV-Vis) to near-infrared (NIR) wavelengths and finite elementfrequency-domain simulation suggest that this MoS₂@TiO₂ heterostructuredphotocatalyst enhances activity for H⁺ reduction. A high H₂ yield rateof 181 mol h⁻¹ cm⁻² (equal to 580 mmol h⁻¹ g⁻¹ based on the loading massof MoS₂) is achieved using a low catalyst loading mass. The spatiallyuniform heterostructure, correlated to plasmon-resonance throughconformal coating MoS₂ that effectively regulated charge transferpathways, is proven to be vitally important for the unique solar energyharvesting and photocatalytic H₂ production.

Examples

Disclosed embodiments of the invention are further illustrated by thefollowing specific Examples, which should not be construed as limitingthe scope or content of this Disclosure in any way.

2.1 Material Preparation for the MoS₂ Embodiment

The anodic growth of TiO₂ nanocavities comprised using titanium foils(25.4×25.4×0.05 mm thick, 99.7% purity, MTI Corporation) as thesubstrates that were ultrasonically cleaned in acetone, ethanol, anddeionized (DI) water for 30 minutes and were then dried in air. Thesample size was rationally chosen but not limited to 1-inch square bythis fabrication method. Anodization was carried out in 3M HF/H₃PO₄(98%, Alfa Aesar, US). A constant voltage of 10 V was applied to atwo-electrode setup for 4 hr with a Pt foil as a counter electrode.After anodization, the TiO₂ films were rinsed with ethanol, then driedin air.

Mo deposition and sulfurization comprised using an e-Beam evaporator(Thermo Scientific) to deposit Mo layers into anodized TiO₂ films with athickness of 10, 20, 30 nm, respectively. The Mo layer thickness wascontrolled by an automated quartz crystal film-thickness monitor. Theas-prepared Mo@TiO₂ hybrid was then put in the center of a quartz-tubefurnace (MTI Corporation) together with 0.5 M Na₂S solution containing 1M sulfur (placed in a crucible at the upstream side, at a temperature of100° C., ramping speed of 2° C. min⁻¹). The MoS₂@TiO₂ heterostructurewas fabricated at 400° C. for 10 min with a heating rate of 5° C.min-under vacuum. The obtained MoS₂ mass loading was estimated bydepositing Mo on compact Ti foils with the same e-beam rate and timementioned above, the thickness of obtained Mo was confirmed by SEM andTEM. Loading of Mo was estimated by multiplying volume and density ofMo. Deposited Mo is supposed to be changed to MoS₂ (10.2, 20.4 and 30.6μg cm⁻² for 10, 20, and 30 nm Mo deposits, respectively).

2.2 Material Characterization

Morphologies of MoS₂@TiO₂ films were observed with a field-emissionscanning electron microscope (FE-SEM, ZEISS ultra 55) and ahigh-resolution transmission electron microscope (FEI Titan 80-300 STEMwith probe Cs corrector and ETEM with imaging Cs corrector). Thecross-section sample for TEM was cut off by a Tescan LYRA-3 Model GMHfocused ion beam microscope and pasted onto Cu ring holder. X-raydiffraction (XRD) was obtained using a X'pert Powder (PANalytical,equipped with a Panalytical X'celerator detector using Cu Kα radiation,λ=1.54056 Å). The chemical composition was characterized by X-rayphotoelectron spectroscopy (XPS, Physical Electronics 5400 ESCA). Ramanspectra were measured using a Renishaw InVia Microscope Raman (532 nmlaser excitation). Absorption spectra were obtained on a Cary WinUV-visible spectrometer from 300 to 700 nm. Incident angle (θ, the angleof excitation light to normal through the center of TiO₂ nanocavity) andexcitation polarization angle (φ, excitation polarization around thenormal through the center of TiO₂ nanocavity) dependent absorptionspectra were tested by rotating the sample from 0-180° with an fixedincident angle or rolling over the sample from a vertical direction to aparallel one. A Fourier transform infrared spectroscope (spectrum 100FT-IR spectrometer, PerkinElmer) was employed for testing absorbance inthe infrared region. The photoluminescence spectra were collected by aNanoLog Spec fluorescence spectrometer excited by a helium-cadmium lampat 400 nm.

Regarding XPS fitting, all spectra were analyzed with the Casa XPSsoftware (version 2.3.15). The samples were conductive so the bindingenergies were not charge referenced. Shirley background subtraction wasused for all spectra. Peak models for each photoelectron line weregenerated using nonlinear least-squares Gaussian/Lorentziancurve-fitting. The S 2p line was fitted with multiple sets of doubletsand constrained to have the same FWHM within each pair as well as a 2:1intensity ratio and 1.16 eV separation between the 2p_(3/2) and 2p_(1/2)peaks. Similarly, the Mo 3d line was fitted with multiple sets ofdoublets as well as a 3:2 intensity ratio and 3.3 eV separation betweenthe 3d₅/2 and 3d3/2 peaks.

2.3 Photocatalysis Characterization

Regarding photocatalytic H₂ evolution, H₂ production was performed in asealed 20 ml quartz reactor. MoS₂@TiO₂ films with a size 7×7 mm weresubmerged in 15 ml of a mixed solution made of DI water (Sea water) andmethanol (8:2 by volume). Subsequently, the reactor was illuminated by asolar light simulator (AM 1.5, 300 W Xe, 100 mW cm⁻²) or monochromaticlight (Zahner CIMPS-QEIPCE system with monochromator and light sourcefrom 350 to 800 nm). The gas produced from the upper space above thesolution in quartz reactor was periodically analyzed.

2.4 Computational Method

Regarding finite Element Method Simulation (FEM), COMSOL MULTIPHYSICS(version 5.2) was used for the FEM solution. The 3D simulation model wasdesigned as a simplified heterostructure, where three MoS₂ nanorods(diameter: 10 nm; length: 20 nm, parallel to the y-axis) standingvertically on TiO₂ (diameter: 50 nm; length: 70 nm, parallel to thez-axis). The incident laser wavelength was set from 300-700, with thepolarization direction along the x-axis. The electric fielddistributions of hybrid nanocavity arrays were monitored across themiddle of TiO₂ nanotube at x-z and y-z planes.

Regarding band structures calculation of MoS₂. First-principlescalculation based on density functional theory (DFT) was used toestimate the electronic property of MoS₂ layer upon S-vacancy deviation(Vienna ab initio simulation program, VASP). The ion-electroninteractions were depicted by projector-augmented wave method. Thegeneralized gradient approximation (GGA) was adopted withPerdew-Burke-Ernzerhof (PBE) exchange-correlation function. 4×4 unitcells were employed as calculation model with S-vacancy of 0%, 8%, and16% respectively. Band structure was relaxed by a 6×6×1 special k-pointmesh grid with 880 eV energy cut off on plane-wave.

3. Results and Discussion 3.1 Morphology, Microstructure, and ComponentAnalysis

A typical fabrication route of a MoS₂@TiO₂ plasmonic heterostructure isschematically illustrated in FIG. 12A. Highly-ordered honeycomb-shapedTiO₂ nanocavities are obtained through Ti anodization (FIG. 12A, with anaverage pore size of 50 nm and a wall thickness of 10 nm). E-beamevaporation was performed to deposit 10, 20, and 30 nm of Mo onto theanodized TiO₂ nanocavity arrays. An obvious cavity wall thickening(increased from 10 nm to 20 nm) is found after Mo coating (FIG. 12B).

CVD sulfurization was carried out at 400° C. for 10 min on TiO₂ coatedwith different Mo thicknesses (abbreviated MoS₂₍₁₀₎@TiO₂, MoS₂₍₂₀₎@TiO₂,MoS₂₍₃₀₎@TiO₂, respectively for 10 nm, 20 nm and 30 nm thick MoS₂,respectively). It was observed that pore size significantly shrunk afterCVD treatment when the deposited Mo thickness increased from 10 to 20and 30 nm (FIGS. 12B-D, FIG. 12C). Transmission electron microscopy(TEM) of MoS₂@TiO₂ nanocavity arrays shows that MoS₂ nanoflakes aregrown inside TiO₂ nanocavities with a highly-ordered 3D laminatestructure (FIGS. 12A-C). These MoS₂ nanoflakes typically consist of lessthan 10 layers, which were verified by Raman spectra. Part of the MoS₂nanoflakes stand vertically on the TiO₂ surface or connect with TiO₂nanocavity walls at a big intersection angle.

High-resolution TEM (HR-TEM) identified lattice fringes of 0.62 nm and0.35 nm, corresponding to the (002) hexagonal facets of MoS₂ and the(101) facets of anatase TiO₂, respectively. It is also revealed that theMoS₂ nanoflakes are bound seamlessly to the TiO₂ nanocavity wallsurfaces, indicating a perpendicular growth of MoS₂ nanoflakes on theTiO₂ surface where TEM images showed perpendicular growth of MoS₂.Spherical aberration corrected high angle annular-dark-field (HAADF)scanning transmission electron microscopy (STEM) and bright-field (BF)STEM were performed to clarify the atomic structure of MoS₂. The imagecontrast of BF-STEM exhibits a relationship with respect to atomicnumber (Z), therefore the sandwich structure of MoS₂ is clearlyobserved. In addition, the S-vacancies can be identified, as marked byblue dashed circles and blue arrows. Interestingly, it seems theS-vacancy exhibits local ‘staging’ structures, as seen for most S atomsthroughout the structural layers.

Cross-sectional STEM images were obtained of MoS₂@TiO₂ heterojunctions,demonstrating that a fully filled MoS₂ laminated network inside the TiO₂nanocavities was successfully fabricated. Energy dispersive X-rayspectrometry (EDS) mapping analysis showed the distribution of eachelement at interface of junction. The separation assignment of Ti to Moand S further proves the isolated vertical connection of MoS₂ and TiO₂at the heterojunction.

In a normal S steam based CVD process, Mo conversion takes place almostspontaneously upon S arriving at the metal surface, making it difficultto control the MoS₂ morphology. One method to allow more morphologicalcontrol is to slow S evaporation using a liquid precursor containingS_(x) ²⁻ ions (see more details in experimental section) and combinedwith a low-temperature ramping rate (2° C. min⁻¹). The S source derivedfrom the Na₂S_(x) solution first arrives at the hollow TiO₂ nanocavitiesand then will transmit gradually along the cavity walls. As it isunderstood, the growth direction of MoS₂ nanoflakes is synchronous withS diffusion, from the middle of TiO₂ cavity to the edges. Eventually, avertically laminated MoS₂@TiO₂ heterostructure was formed inside theTiO₂ cavities. A photograph of the MoS₂₍₁₀₎@TiO₂ heterostructureprovided in FIG. 12G shows uniform morphologies at different samplelocations, indicating a highly uniform MoS₂@TiO₂ throughout the entiresample. Control experiments were carried out by using a normal solid Ssource, flat TiO₂ film or Mo foil in duplicate reactions. Larger MoS₂flakes with disordered architecture were obtained for all controlexperiments (FIGS. 12D-F), further verifying the significant impact madeby the addition of TiO₂ nanocavity arrays and using liquid S sourcetoward the inhibited MoS₂ overgrowth that promotes the formation of aMoS₂@TiO₂ heterostructure configuration.

XRD patterns of as-prepared MoS₂@TiO₂ films were obtained. The maindiffraction peaks are indexed to the Ti substrate (PDF file No. 44-1294)because of the thin thickness of the TiO₂ film and low MoS₂ content.Noticeably, a weak peak at 14.5°, corresponding to the c-plane (002) of2H—MoS₂ (S—Mo—S graphene-like layer, PDF file No. 86-2308), can beobserved and intensifies with Mo thickening. Diffraction patterns of Tisubstrates are a little different by the orientation diversity ofcommercial Ti foil, depending on the manufacturing process. Ramanspectroscopy is more distinct for demonstrating the MoS₂ layeredstructure. Peaks around 378 and 402 cm⁻¹, corresponding to in-planeE_(2g) and out-of-plane A_(1g) vibration modes of 2H—MoS₂ are dominantin the spectra. Red-shift of E_(2g) and blue-shift of A_(1g) occurs withincreasing Mo thickness from 10 to 30 nm as well as prolongingsulfurization time from 10 to 50 min for the MoS₂₍₃₀₎@TiO₂ film.

A frequency difference monotonically increased from 21.25 cm⁻¹ to 24.24cm⁻¹, which is in excellent agreement with the literature reportindicating a thickness of less than 10 stacking layers of MoS₂ perflake. The greater intensity ratio of A_(1g) to E_(2g) is due to a largeamount of surface edge exposure and strong interlayer restoring forceinteractions acting on misalignments or defects in the MoS₂ layers. Itis very clear that there is no Mo oxide can be found from both XRD andRaman, indicating a complete conversion from Mo to MoS₂ after CVDtreatment. Experimental tests also indicate the effect of Mo oxide onthe optical and photocatalytic properties can be ignored, which will befurther discussed in the following sections.

Surface defects, chemical states and the stoichiometric ratio of Mo andS in the MoS₂@TiO₂ heterostructure were investigated by X-rayphotoelectron spectroscopy (XPS) analysis. 2H—MoS₂ and TiO₂ wereconfirmed by XPS results). The high-resolution S 2p spectrumdemonstrates that there are three overlapping chemical states of S. Thedoublet centered at 161.7 eV is assigned to S 2p in a Mo—Sconfiguration, the upshifted peaks at 163.4, 166.6 and 169.9 eV are dueto the existence of S, Mo—O—S band or oxidized S on the surface of MoS₂.More oxidized S was found with an increase the amount of MoS₂. Bindingenergies for Mo 3d_(3/2), Mo 3d_(5/2) are fitted to a pair of doubletsat 227.9/231.2 eV and 228.9/232.2 eV, respectively, which confirms theMo³⁺ and Mo⁴⁺ accordingly. The existence of Mo³⁺ ions implies theformation of S defects in MoS₂ and account for about 20-30% of the totalMo element semi-quantitatively estimated from the XPS spectral area inthe MoS₂₍₁₀₎@TiO₂ sample (S-vacancy is in a range of 5-8% percentage ofS atoms deficiency). Mo⁶⁺ ions can be slightly detected because ofsurface oxidation during in-air storage. Combining the aforementionedinvestigations, it can be concluded that the S-vacancy andnonstoichiometric features of MoS₂ nanoflakes are where free electronsand plasmonic resonance processes originate. The XPS spectral states ofTi and O are in agreement with crystalline TiO₂.

3.2 Spectral Response and Photocatalytic Performance

The solar light harvesting of MoS₂@TiO₂ heterostructure was investigatedthrough UV-Vis absorption spectra with an incident angle of 0°. Theintrinsic absorption edge located at about 360 nm belongs to the TiO₂substrate, suggesting a band gap of 3.2 eV. Furthermore, very broadpeaks ranging from 400 to 600 nm can be detected on the MoS₂@TiO₂heterostructures. This visible light response is distinctly far fromband gap excitation of both MoS₂ and TiO₂ portions, consistent withheterostructured films. Consequently, a metal-like LSPR is assigned tothis absorbance, which arises from collective oscillations of excesscharges (electrons) on the edge of MoS₂. Abundant S-vacancies and highlyordered vertically laminated structures may dominate the free chargeinteractions. It has been well recognized that the catalytic activitiesof both semiconducting 2H—MoS₂ and metallic 1T-MoS₂ generally arise fromedge sites and S-vacancies. Both edges and S-vacancies are consideredresponsible for crystal asymmetry of the analyte, where the electronicstructure is changed slightly. Electrons in a higher Fermi level areeasy to transfer to local collected vacancies on the surface of MoS₂.When the orbital density vibration of these electrons resonant withincident light, photo-excited local dipoles and charge separationhappens, forming a confined local field around MoS₂ surface. TherebyS-vacancies and surface “hot” electrons are applied to plasmonicantennas, which arises a modulation of the Fermi level and drasticallychange the initial electronic properties of MoS₂.

Combining the plasmonic effect with the intrinsic absorption ofMoS₂@TiO₂, a Full-solar-spectrum harvesting is nearly achieved in thisco-catalyst heterostructure. Additionally, the plasmonic resonance isdemonstrated to be tunable upon Mo mass loading changes from 10 to 30nm, with absorption peaks at 420, 480 and 510 nm, accordingly. Thisprocess may be similar to the classic size-dependent plasmonic effect.The maximal light harvesting cross-section is obtained in the samplewith 30 nm of Mo loading. Control experiments carried out on puremetal-based film (30 nm Mo on TiO₂) and oxidized sample (MoO₃ on TiO₂)give little visible light coverage, but the UV portion from TiO₂substrate. Incident angle and excitation polarization angle-dependentabsorption spectra were also tested on MoS₂₍₃₀₎@TiO₂ (see FIG. 13A). Theresonance wavelengths are independent of the polarization angle (seeFIG. 13A). All the absorption peaks of MoS₂₍₃₀₎@TiO₂ is located in arange of 450-500 nm at an incident angle of 15°. The absorbancefluctuation (cross-section of curves) may come from the nonuniformity ofMoS₂ flakes inside TiO₂ (as revealed by TEM). Absorption spectrameasured at different incident angles were present. The absorption peakposition increases from 490 nm to 600 nm with the increase of 0 to 90°.

The absorption peak position can further shift backward when theincident angle decreases to 0°. This dynamic shifting indicates that theresonance has 90° difference in the incident angle dependence,suggesting that the orientations of MoS₂ flakes affect the resonancemodes and charge distribution profiles in the hybrid catalyst. Theexcitation efficiencies of the samples are also dependent on incidentangles. A maximum photo-efficiency (absorption cross-section) isachieved at an incident angle of 30°, consistent with the previousfindings that the resonance intensity can be optimized at certainangles. The relationship of nonstoichiometric features to lightharvesting was verified by extending the sulfurization time from 10 minto 50 min on MoS₂₍₃₀₎@TiO₂ film (see FIG. 13B). A strong absorbancedegeneration can be observed with longer reaction times andcomplementary to the S source within the heterostructure, implying thatS-vacancy is the main plasmonic charge donor.

Surface-enhanced Raman scattering (SERS) was used to further determinethe enhanced electromagnetic field in the vicinity of MoS₂@TiO₂.Mercaptobenzoic acid (4-MBA) was used as the probe molecule with a laserexcitation at 630 nm. The dominant peaks found to be at 1099 and 1594cm⁻¹ are attributed to the vibration mode of v8a aromatic ring andbreathing mode of v12a ring, respectively. The weak peaks at 1293 and1188 cm⁻¹ are the stretching mode of ν_((COO—)) and deformation mode ofC—H, respectively. The stronger SERS signal on a MoS₂@TiO₂ substrate(comparing with bare TiO₂ nanocavity arrays using 10⁻⁵ mol L⁻¹ 4-MBA)demonstrates the plasmonic polarization presented on the MoS₂@TiO₂surface.

Photocatalytic activity of series MoS₂@TiO₂ co-catalysts for H₂evolution was estimated under simulated solar light illumination withmethanol as a hole scavenger in DI water. Pure TiO₂, sulfurized TiO₂,and MoO₃@TiO₂ hybrid are photocatalytically inert toward hydrogenevolution, with a reaction rate of lower than 10 μmol h⁻¹. Noticeably,the photocatalytic activity is sharply enhanced with MoS₂ loading andachieves maximal values on the MoS₂₍₃₀₎@TiO₂, reaching 84 μmol h⁻¹(≈8times that of pure TiO₂). The photocatalytic activities of the samplesincrease non-linearly with increasing of MoS₂ loading (from 10 to 30 nmMo deposition) because of the relative lower solar energy input viaincreasing catalysts amount. Additionally, a deterioration is found withlonger sulfurization time from 10 to 50 min (MoS₂₍₃₀₎@TiO₂), deducing aweakened plasmonic effect by the stoichiometric assignment, according tospectroscopy analysis (see FIG. 13B). Photocatalytic activities of thesamples were obtained, where an optimal photocatalytic activity of 181μmol h⁻¹ cm⁻² is observed on the heterostructure coated with 30 nm Mo.FIG. 14A shows the rate of H₂ production normalized by mass loading ofMoS₂ (mass-measurement of MoS₂ catalyst is given in the experimentsection). Samples with 10.2 μg MoS₂ has an activity of 580 mmol h⁻¹ g⁻¹,which is superior to the state-of-the-art photocatalytic systems shownin Table 1 below.

TABLE 1 Table 1. Photocatalytic performance comparison of disclosedMoS₂/TiO₂ based catalyst for H₂ evolution (bottom table entry) to otherstructures. H₂ evolution rate Mass activity Method Materials (μmol h⁻¹)(μmol h⁻¹g⁻¹) Solvothermal MoS₂/C₃N₄ nanograss 55.6    11.1 CVDMoS₂/TiO₂ nanosheets 42  4,200 Hydrothermal MoS₂/TiO₂ nanorods  2.7 1,500 CVD MoS₂/TiO₂ nanotubes  0.44 — CVD MoS₂/TiO₂ 89 580,000nanocavities

That may be ascribed to the highly-ordered architecture of MoS₂nanoflakes on TiO₂ nanocavities with a lower mass loading, which reducesthe electron-hole recombination. Furthermore, LSPR induced electronenergetic level is more negative in the MoS₂₍₁₀₎@TiO₂ because of theblue-shifted light harvesting and more negative lowest unoccupiedmolecular orbital level that favors effective charge utilization ratesin H⁺ reduction. Stability and recyclability of the MoS₂₍₁₀₎@TiO₂co-catalyst were estimated by repeating intermittent H₂ evolution undersimulated solar light. 90% of incipient production can be kept for thehybrid films even after 21 h. Control experiments were carried out onthe samples using MoS₂ deposited on compact TiO₂ and Mo foils(MoS₂₍₃₀₎@TiO_(2(compact)) and MoS₂@Mo_((compact))) to further confirmthe crucial role of TiO₂ nanocavity arrays in the improvedphotocatalytic stability and recyclability. We found that MoS₂ flakesfell off from these compact substrates during H₂ evolutions testing,leading to a considerably deteriorated H₂ evolution. As shown in FIG.14B, C, approximately 30% of H₂ yield (less than 15 μmol h⁻¹) was keptfor the MoS₂@TiO_(2(compact)) after 5 circles testing. An even less H₂yield (about 20%, equal to 10 μmol h⁻¹) was observed on theMoS₂@Mo_((compact)) after 15 h testing. These control experimentsindicate that TiO₂ nanocavity arrays not only contribute to UV-lightabsorption but also serve as host to immobilize MoS₂ co-catalyst.

Photoluminescence (PL) spectrum is used to detect the charge carriertrapping, migration and transfer mechanism during photocatalyticreactions (FIG. 14D, E). The steady-state PL spectra located at 595 nm(FIG. 14D), characterizes remarkable quenching along with the Mothickening from 10 nm to 30 nm, suggesting either a shorter lifetime orfaster trapping of photo-electrons with increasing MoS₂ content.Transient-state PL decay curves of MoS₂@TiO₂ heterostructures arecompared in FIG. 14C. A bi-exponential function was used here forfitting the PL decay curves mathematically:

y=A ₁exp(−x/τ ₁)+A ₂exp(−x/τ ₂)+y ₀  (1)

where A₁, A₂, and y₀ are amplitude coefficient and basal constant. τ₁and τ₂ are fluorescent lifetimes corresponding to non-radiativerecombination and inter-band recombination, respectively. The calculatedcarrier lifetime is inserted in FIG. 14E accordingly. The averagefluorescent lifetime of MoS₂₍₁₀₎@TiO₂ (τ₁: 2.8 ns and τ₂: 14 ns) islonger than the other ones after integrating by the equation:

$\begin{matrix}{\mspace{79mu}{{\tau = \frac{\text{?} + \text{?}}{\text{?} + \text{?}}}{\text{?}\text{indicates text missing or illegible when filed}}}} & (2)\end{matrix}$

This indicates that appropriate MoS₂ loading and periodically patternedmorphology provide the MoS₂₍₁₀₎@TiO₂ film with less trapping centers andprolonged retention of hot electrons, as well as enhanced H₂ production.

3.3 Computational Study

The photo-electron arising mechanism, transfer pathway and consumingprofile were computationally studied to assist in explaining theexcellent photocatalytic activity of MoS₂@TiO₂ heterostructures.Firstly, the UV portion of the solar light is mainly absorbed by theTiO₂ substrate, where photo-excited electrons transfer to the MoS₂ basalplane and diffuse to edge active sites. Here, the designed MoS₂@TiO₂short nanocavity arrays and seamless junctional connection dramaticallyreduce the charge barriers at active sites. The perpendicular MoS₂—TiO₂configuration benefits the electron transfer pathways at the basal planewhere a lower resistance and suppressed quenching capture are expected.On the other hand, a broad LSPR band (ranging from 400 to 600 nm)renders the MoS₂@TiO₂ heterostructures for visible light-driven H⁺reduction. This plasmon-enhanced activity is presumably a result of theaforementioned S-vacancies (confirmed by the STEM) and providespredominant aspects for H₂ evolution. Finally, photoelectrons excited onthe MoS₂ conduction band from NIR light can also contribute to H₂evolution. Although, considerable carrier recombination is unneglectablein this narrow band gap semiconductor. MoS₂ edge defects (S-vacancies)play a dominant role in plasmonic resonance and hot electron excitationupon solar light irradiation, but will also capture free carriers andcause unwanted quenching. Moreover, MoS₂ and TiO₂ intrinsicrecombination is another way for hot electron consumption in catalyticprocesses (dash line between band gap), which may weaken photocatalyticperformance.

To have an insightful understanding of the inter-band excitation of MoS₂nanoflakes and eliminate the possibility of MoS₂ intrinsic activity invisible-light, electronic band structure was calculated by densityfunctional theory. S-depletion was introduced in the computation modelswith an atom content of 0%, 8%, 16% respectively, mainly at the edgeposition. A fundamental band gap around 2.07 eV is found here for aperfect MoS₂ crystal, with the highest occupied valence band and thelowest empty conduction band located at 1.63 eV and −0.4 eV,accordingly. The band structure is modulated by introduction of anS-vacancy, which induces the narrowing of the band gap to 1.73 (8%) and1.12 eV (16%). Moreover, defect states appear near to the Fermi levelfor the model with 16% S-depletion, which improves the metallic featureand charge mobility of S-depleted MoS₂ layers. The overall electronicdensity of states (DOS) for MoS₂ before and after S-depletion wereobtained. Defect states near to the Fermi level originate mainly fromthe d orbitals of Mo and p orbital of S atoms, creating pseudo-ballisticelectron transport channels within MoS₂. This result is furthersupported by other literature and fully compatible with optical spectrain the IR region), where the MoS₂ band gap absorption peaks are locatedaround 882 nm (1.40 eV). A further conclusion can be deduced from thediscussion aforementioned that solar light caused activity of MoS₂@TiO₂films mainly functionalized by the plasmonic effect surrounding MoS₂surface.

To further confirm this assumption, wavelength-dependent photocatalysisand finite element method (FEM) simulations were performed on MoS₂@TiO₂heterostructures under monochromatic light illumination (350-700 nm,with a 50 nm interval). A spectrum dependent H₂ evolution rate isrevealed over the UV-Vis-NIR region, with a consistent variation inabsorption spectra for the MoS₂@TiO₂. The largest H₂ yield rate is about59, 73, 86 μmol h⁻¹ mW⁻¹ cm⁻² for Mo loading of 10, 20 and 30 nm,respectively, located at 450, 500, 550 nm accordingly. The relativeconsistency of photo-absorption and H₂ production over solar lightindicates that the photocatalytic activity enhancement is primarilydriven by the plasmonic effect from blue-green wavelengths. Uponexcitation by monochromatic light irradiation, the H₂ yield is foundhigher than equivalent solar light (AM 1.5 100 mW cm⁻²) excitation,especially in the 400-600 nm region. This suggests an additional benefitfor use of variant colored light for plasmonic material based H₂generation. A 3D model was designed using a single TiO₂ nanocavity (50nm in diameter and 70 nm in length) and three vertically loaded MoS₂nanorods (10 nm in diameter and 20 nm in length) on both sides. It wasfound that light irradiation coupled with LSPR yields strong electricfield enhancements at the tip of MoS₂. The field intensity increaseswith wavelength from 300 to 400 nm and reduces significantly after 500nm, correlating the LSPR absorption spectra of MoS₂₍₁₀₎@TiO₂ film. Themaximum electric field enhancement is produced at 400 nm illumination.Field intensity between each antinode covers a distance longer than 5nm, achieving zero-gap field distribution between MoS₂ interspace (<10nm). Consequently, a 3D electric field distribution is built throughoutthe networked MoS₂ laminate heterostructure and responds strongly tophotocatalytic H₂ production.

Prospectively, H₂ generation from seawater will be highly desirable.However, few studies have focused on the practical photocatalysis onnatural seawater because of the barrier blocking of dissolved salt forphotocatalytic activity and durability of catalysts. Here it wasinvestigated the effect of salt (mainly NaCl) on photocatalytic activityof MoS₂@TiO₂ films with natural seawater (pH: 8.4) splitting under thesimulated solar light. Obviously, the activity from seawater decreasedmarkedly compared to that of pure water. Because the isoelectric pointof MoS₂ is lower than 7 by the plasmonic “hot” electrons on the surface.Electrostatic adherence of hydroxyl group, metal ions, and oxidation ofsulfides take place severely in the seawater with a weak alkalineenvironment. Fortunately, 60% of incipient production is kept forsamples, where the highly-ordered architecture may be helpful. Thiscontrol experiment permits our MoS₂@TiO₂ heterostructure as a promisingmaterial for H₂ production from seawater.

While various disclosed embodiments have been described above, it shouldbe understood that they have been presented by way of example only, andnot as a limitation. Numerous changes to the disclosed embodiments canbe made in accordance with the Disclosure herein without departing fromthe spirit or scope of this Disclosure. Thus, the breadth and scope ofthis Disclosure should not be limited by any of the above-describedembodiments. Rather, the scope of this Disclosure should be defined inaccordance with the following claims and their equivalents.

1. A heterostructured catalyst, comprising: a 2-dimensional (2D) array of titanium comprising nanocavities that are all directly attached to a substrate, wherein each of the titanium comprising nanocavities have a pore with a nanopore size and a wall with a nanowall thickness.
 2. The heterostructured catalyst of claim 1, wherein the titanium comprising nanocavities comprise titania nanocavities, further comprising a metal layer or a metal compound layer on the titania nanocavities including inside the pores.
 3. The heterostructured catalyst of claim 2, wherein the titanium comprising nanocavities consist of SrTiO₃ including a surface layer consisting of reduced SrTiO₃.
 4. The heterostructured catalyst of claim 2, wherein the 2D array of the titanium comprising nanocavities is a periodically ordered array having an interval of a periodicity of 10 nm to 100 nm with a distribution all within ±1 nm of a mean interval value, the nanopore size is 10 to 200 nm with a distribution all with ±1 nm of a mean pore size value, and the nanowall thickness is 5 nm to 20 nm with a distribution all within ±0.5 nm of a mean wall thickness value. 